Process for making strain-hardened polymer products

ABSTRACT

A process which subjects a body of polymer to deformation to produce strain hardened polymeric products including blending a polymer and a nanoparticle material to produce a polymeric composition, forming a film from the polymeric composition, and subjecting the composition to strain hardening. The resulting product has improved clarity, dimensional stability, uniform thickness, due to the “self-leveling” properties of the polymeric composition.

BACKGROUND OF THE INVENTION

The present invention is related to a process for making strain-hardenedpolymer products using polymer-nanoparticle compositions.

Strain hardening is the most important characteristic in the productionand usability of polymeric films in such processes as Tenter framebiaxial film stretching, blow molding, film blowing, thermoforming, andthe like. This property affects structural characteristics such ascrystallization as well as thickness uniformity of polymers. It is atthis point that many polymers are known to crystallize due toorientation from deformation, with a consequential affect on optical,physical and mechanical properties.

The reasoning for controlling strain hardening is due to its effect onstructural development in materials. The net effect of this behaviortranslates into “self-leveling” that allows thinned portions of amaterial to sustain higher forces while transferring the deformation toother undeformed regions, resulting in achieving greater uniformity inthe thickness of the products. For this reason, it is of criticalimportance to find an efficient means of selectively controlling thestrain hardening behavior of polymers. As an example, in film productionfor such industries as information recording (audio and video cassettes,etc.), film uniformity and lack of surface roughness are of paramountimportance for acceptable products.

There have been many attempts to control the strain hardening behaviorof polymeric films through the addition of additives, by blending, or byco-polymerization. But, adding fillers does not necessarily improvemechanical and physical properties of polymers, and studies indicatethat they may not have desired effect on the deformation behavior. Forexample, the work of Taniguchi et al (Atsushi Taniguchi and Miko Cakmak,“The effect of titanium dioxide particles on the deformation behaviorand orientation development in PET films”, Annual TechnicalConference—Society of Plastics Engineers (2000), 58th (Vol. 2),1786-1790) showed that adding submicron sized particles retards thestrain hardening process to higher strain levels. Tanaguchi et al.reported that the effect of submicron size TiO₂ particles at varyingconcentrations on the stress-strain behavior of uniaxially deformed PETfilms from the amorphous state. The TiO₂ particles act as nucleationagents and enhance the thermally induced crystallization of PET. Whenstretched from the amorphous state, TiO₂ particles at concentrations aslow as 0.35 percent reduce the overall stress and delay strainhardening, thereby hindering orientation induced crystallization.Consequently, films stretched under the same conditions, but containinghigher levels of TiO₂ have both lower crystallinity and orientation.They attribute this behavior to the reductions in the number of chainentanglements due to the presence of small amounts of TiO₂ particles.

Iwakura, et al. (Iwakura, K.; Wang, Y. D.; and Cakmak, M., “Effect ofbiaxial stretching on thickness uniformity and surface roughness of PETand PPS films,” Int. Polym. Process. (1992), 7(4), 327-333) in theirpublication performed biaxial film-stretching studies with poly(ethyleneterephthalate) (PET) and poly(p-phenylene sulfide) (PPS) to observe theeffects of stretching conditions on film properties; particularlysurface smoothness and thickness uniformity. They found that bydecreasing the stretching temperature and increasing the stretch ratio,they could improve these properties in PET, but not in PPS. They alsoobserved that the most important factor in controlling these propertieswas the strain hardening mechanism. If this could be controlled, so toowould the thickness uniformity and surface smoothness. Under theconditions used in these experiments, strain hardening occurred for PET,but not for PPS, and once strain hardening was attained, propertiesimproved drastically, especially the thickness uniformity. Theyattributed the problems in PPS to branching of the polymer chains.

It would be a very easy to process films to have desirable properties ifall that one had to do was find the strain hardening point, and juststretch beyond this. However, this is not always possible. Sometimes thestrain hardening point occurs at too high of a strain to be feasible, orin some other cases, strain hardening is just not possible under normalprocessing conditions. In addition, there are other side issues inprocessing and properties. Characteristics, such as miscibility anddomain size, have great effect on polymer properties. In the instancewhere strain hardening can be improved, other properties may suffer. Amodification that works very well in one case can yield a very poorresult in another, or too much modification may actually be detrimentalto some properties. In other cases, a modification may work very well,but the strain hardening behavior still might not be very wellcontrolled, as the modification might work only to a certain extent.

SUMMARY OF INVENTION

The present invention is the result of the discovery that a way tocontrol the strain hardening behavior of polymer films during stretchingfrom their rubbery state temperature range from their amorphousprecursors through the inclusion of small fraction of nano sizedparticles. This discovery specifically provides a new highlycontrollable “tool” to achieve a deformation behavior of the polymericmaterials, wherein the strain hardening behavior can be adjusted tooccur at different desired strain levels thereby causing the films to“self level” for different product geometries without having to changethe processing temperature and processing stretching rates.

There are also additional advantages to these polymer-nano particlecomposites. Among these include increased heat distortion temperature,modulus, gas barrier properties. In addition at the small concentrationlevels used in the composition, the films remain transparent in thevisible wavelength range. The latter is very important in applicationswhere the product appearance is important (e.g., beverage containers,transparencies, etc.).

It is therefore, an object of the present invention to provide a methodto control the strain hardening behavior of polymer films usingnanometer scale particles.

At least one or more of the foregoing objects, together with theadvantages thereof over the known art relating to polymer films, whichshall become apparent from the specification which follows, areaccomplished by the invention as hereinafter described and claimed

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a graph of the True Stress versus True Strain ofcompression-molded poly(L-lactic acid) films having Montmorillonite clayloadings of 0, 1 and 3% by volume and stretched at 75° C. and at 5% perminute.

FIG. 2 is a graph of the True Stress versus True Strain ofcompression-molded poly(L-lactic acid) films having Montmorillonite clayloadings of 0, 1, 3 and 10% by volume and stretched at 75° C. and at 50%per minute, then held in constrained length at 75° C. for 30 minutes.

FIG. 3 is a graph of the True Stress versus True Strain ofcompression-molded poly(L-lactic acid) films having Montmorillonite clayloadings of 0, 1, 3 and 5% by volume and stretched at 75° C. and at 500%per minute.

FIG. 4 is a graph of the True Stress versus True Strain ofcompression-molded poly(L-lactic acid) films having Montmorillonite clayloadings of 0, 1 and 3% by volume and stretched at 85° C. and at 500%per minute.

FIG. 5 is a graph of the True Stress versus True Strain ofcompression-molded poly(L-lactic acid) films having Montmorillonite clayloadings of 0, 1 and 3% by volume and stretched at 95° C. and at 500%per minute.

FIG. 6 is a graph of the True Stress versus True Strain ofextrusion-cast poly(L-lactic acid) films having Montmorillonite clayloadings of 0 and 3% by volume and stretched at 75° C. and at 5% perminute.

FIG. 7 is a graph of the True Stress versus True Strain ofextrusion-cast poly(L-lactic acid) films having Montmorillonite clayloadings of 0 and 3% by volume and stretched at 75° C. and at 50% perminute.

FIG. 8 is a graph of the True Stress versus True Strain ofcompression-molded poly(L-lactic acid) films having Montmorillonite clayloadings of 0 and 3% by volume and stretched at 75° C. and at 500% perminute.

FIG. 9 is a graph of the True Stress versus True Strain ofextrusion-cast poly(L-lactic acid) films having Montmorillonite clayloadings of 0 and 3% by volume and stretched at 85° C. and at 500% perminute.

FIG. 10 is a graph of the True Stress versus True Strain ofextrusion-cast poly(L-lactic acid) films having Montmorillonite clayloadings of 0 and 3% by volume and stretched at 95° C. and at 500% perminute.

FIG. 11 is a graph of the True Stress versus True Strain ofcompression-molded polystyrene films having Montmorillonite clayloadings of 0, 1 and 3% by volume and stretched at 115° C. and at 1000%per minute.

FIGS. 12A and 12B are schematic representations of an as-cast(underformed) nanocomposite film and a deformed (axially drawn)nanocomposite film.

DETAILED DESCRIPTION AND PREFERRED EMBODIMENT FOR CARRYING OUT THEINVENTION

The present invention is directed to a process which subjects a body ofpolymer to deformation to produce strain hardened polymeric products. Byadding nanoparticles to the polymers, when the polymer composition issubjected to deformation, novel strain hardening characteristics areimparted to cast films. This is applicable to a wide variety of filmforming materials including, but not limited, to homopolymers andcopolymers of polyolefins, polyamides, polyimides, polyesters, aliphaticpolymers, amorphous polymers, slow crystallizing polymers, fastcrystallizing polymers, and blends, alloys and combinations of the same.Depending upon the polymer, the effect on strain hardening is alwayspresent, but in varying degrees. In fast crystallizing polymer thestrain hardening can be achieved when the polymer is either partiallymolten or completely molten.

Nanoparticles would include particles of a nanoscale, i.e., 400 to 700nanometers, with at least one dimension in the nanoscale, includingspheres, particles of irregular geometry, sheets and foils, and fibers,wires and tubes, such as:

-   -   Carbon nanoparticles (graphite, nanotubes, spherical particles        such as Bucky Balls, and the like)    -   Glassy nanoparticles, including but not limited to silica-based        nanoparticles    -   All varieties of nanoclays, with substituted Montmorillonite        being preferred    -   Metal oxide, metal sulfides, metal nitrides, and other such        ionic nanoparticles    -   Metal complex nanoparticles    -   Metallic and metallic alloy nanoparticles (nanowires,        nanospheres, nano-sized sheets and foils)    -   Colloidal nanoparticles

The amount of nanoparticles will vary depending upon the polymeremployed and the characteristics of the polymer product, but will varybetween 0.01 percent by volume up to 10 percent by volume, with therange of 0.1 percent to 10 percent by volume nanoparticle loading beingpreferred, and the range of 1 percent to 10 percent by volume being alsopreferred. Usually the least amount of nanoparticle to be effective willbe selected, since the additional material only increases materialcosts. But, there may be instances where higher loadings are desirabledue to the final properties of the final polymeric product. Further,since at increased loadings, the strain hardening point continues toshift to lower true strains, by varying the nanoparticle concentrationone can “dial-in” the strain hardening point for a specific process.

The polymeric products can be two dimensional, such as films, or threedimensional, such as bottles or other shapes.

To produce the polymeric products, the polymer and nanoparticles areblended together to form a nanocomposite which is subsequently furthertreated. But, there are several means of making the nanocomposite,including swelling followed by polymerization (i.e., mixing thenanoparticles with the chemicals before the polymer is made), rubberystate blending, melt compounding (where the polymer is completelymelted), and the like. This part of the process is not critical. Mixingcan be performed in chemical reactors (in polymerization for example),internal mixers, continuous mixers, calendars, extruders with variousscrew types, configurations, and various numbers of screws, includingsingle, twin, and/or multiscrew mixers, extruders incorporating mixingelements in their screw designs, and any other similar mixinginstruments. The operating conditions vary depending upon the polymerand/or nanoparticles employed. Temperatures range from very low,including below zero in some cases of polymerization, up to very high,including more than 400 C in the case of some high temperature polymerswhere blending occurs in the melt.

When the polymeric materials are used to make films, they can beextrusion-cast, compression-molded or blown. Thereafter, the polymersare quenched to low temperatures (below the glass transition forpolymers that can be quenched into the amorphous state), then thetemperature is raised to a desired stretching temperature. Under theseconditions the polymer can be stretched to yield high orientation andstrength.

Thus, strain-hardened or strain-hardening would include or apply toproducts subjected to deformation in order to secure the specifiedphysical characteristics such as orientation and crystallinity. Thiscould also be defined as the point at which there is a sudden upturn inthe true stress-true strain curve for the composition. This isillustrated in the figures which are part of the present disclosure.These processes include, but are not limited to thermoforming, bottleblowing, stretch bottle blowing, film casting, uniaxial and multi-axialstretching, and various other stretching or film blowing processes.

When the polymeric products are made from a slow crystallizing polymer,such as poly(lactic acid) (PLA), polyetheylene terephthalate (PET),poly(p-phenylene sulfide) (PPS), or poly(ethylene-2,6-naphthalate)(PEN), during film casting, the polymer exits the die as a molten sheetand is quenched below it's glass transition, so that the materialpossesses little or no crystallinity. Then, the temperature of thematerial is raised to somewhere between the glass transition and coldcrystallization temperature, which is called the rubbery state, and thepolymer is stretched (uniaxially or multi-axially), so that the chainsbecome oriented in the direction, due to the stretching. The polymer maystrain harden, or strain induced crystallization will occur, and thismay lead to strain hardening. Amorphous materials, including polystyreneand PMMA, remain essentially amorphous under all processing conditions,and strain harden due primarily to orientation. Highly crystallinepolymers, including polyethylene and polypropylene, are highlycrystalline under all processing conditions, so the nature of theirdeformation temperature is a little different. If they are melted, theywill not hold up to stretching, so they are only partially melted. Theremaining crystalline portion holds the material together while it isstretched, and the bulk of the deformation takes place in the moltenamorphous portion. In each case strain hardening is important incontrolling both the thickness and the uniformity of the thickness.

By controlling strain hardening through modifying the material, via theprocess of the present invention, the need for expensive and complicatedmodifications to capital equipment, or even the purchase of new capitalequipment is eliminated, while the properties of the material areenhanced. This can result in the saving of a considerable amount ofmoney.

EXAMPLES Example 1 Atactic Polystyrene

Atactic polystyrene pellets were exposed to blending in a Brabendarmixer for ten minutes using banbury blades. Blending occurred at 200° C.with a rotor speed of 70 rpm. After the blending was complete, the meltwas quenched in room temperature water, and was subsequently dried undervacuum. Film samples were compression molded at 200° C. for fiveminutes, and were then quenched in room temperature water. The resultingfilm was amorphous.

Example 2 Atactic Polystyrene with 1% by Volume SubstitutedMontmorillonite

Atactic polystyrene pellets were exposed to blending in a Brabendarmixer for two minutes using banbury rotors. Blending occurred at 200° C.with a rotor speed of 70 rpm. After two minutes had passed, 1% by volumeof substituted Montmorillonite (Nanocor I34) was added to the mixer, andthe polymer and clay were blended for eight minutes. After the blendingwas complete, the melt was quenched in room temperature water, and wassubsequently dried under vacuum. Film samples were compression molded at200° C. for five minutes, and were then quenched in room temperaturewater. The resulting film was amorphous.

Example 3 Atactic Polystyrene with 3% by Volume SubstitutedMontmorillonite

Atactic polystyrene pellets were exposed to blending in a Brabendarmixer for two minutes using sigma blades. Blending occurred at 200° C.with a rotor speed of 70 rpm. After two minutes had passed, 3% by volumeof substituted Montmorillonite (Nanocor I34) was added to the mixer, andthe polymer and clay were blended for eight minutes. After the blendingwas complete, the melt was quenched in room temperature water, and wassubsequently dried under vacuum. Film samples were compression molded at200° C. for five minutes, and were then quenched in room temperaturewater. The resulting film was amorphous.

Example 4 Polylactic Acid

Polylactic acid (Natureworks 4041D) pellets were exposed to blending ina counter-rotating twin screw extruder with two shearing elements andone backflow element. Blending occurred at 190° C. with a rotor speed of190 rpm. The melt was quenched upon exit from the die in roomtemperature water, and was subsequently dried under vacuum. Film sampleswere compression molded at 190° C. for five minutes, and were thenquenched in room temperature water. The resulting film was essentiallyamorphous.

Example 5 Polylactic Acid with 1% by Volume Substituted Montmorillonite

Polylactic acid (Natureworks 4041D) was tumble mixed with 1% by volumeof substituted Montmorillonite (Nanocor I34). The dry mixture wasexposed to blending in a counter-rotating twin screw extruder with twoshearing elements and one backflow element. Blending occurred at 190° C.with a rotor speed of 190 rpm. The melt was quenched upon exit from thedie in room temperature water, and was subsequently dried under vacuum.Film samples were compression molded at 190° C. for five minutes, andwere then quenched in room temperature water. The resulting film wasessentially amorphous.

Example 6 Polylactic Acid with 3% by Volume Substituted Montmorillonite

Polylactic acid (Natureworks 4041D) was tumble mixed with 3% by volumeof substituted Montmorillonite (Nanocor I34). The dry mixture wasexposed to blending in a counter-rotating twin screw extruder with twoshearing elements and one backflow element. Blending occurred at 190° C.with a rotor speed of 190 rpm. The melt was quenched upon exit from thedie in room temperature water, and was subsequently dried under vacuum.Film samples were compression molded at 190° C. for five minutes, andwere then quenched in room temperature water. The resulting film wasessentially amorphous.

Example 7 Polylactic Acid with 5% by Volume Substituted Montmorillonite

Polylactic acid (Natureworks 4041D) was tumble mixed with 5% by volumeof substituted Montmorillonite (Nanocor I34). The dry mixture wasexposed to blending in a counter-rotating twin screw extruder with twoshearing elements and one backflow element. Blending occurred at 190° C.with a rotor speed of 190 rpm. The melt was quenched upon exit from thedie in room temperature water, and was subsequently dried under vacuum.Film samples were compression molded at 190° C. for five minutes, andwere then quenched in room temperature water. The resulting film wasessentially amorphous.

Example 8 Polylactic Acid with 10% by Volume Substituted Montmorillonite

Polylactic acid (Natureworks 4041D) was tumble mixed with 10% by volumeof substituted Montmorillonite (Nanocor I34). The dry mixture wasexposed to blending in a counter-rotating twin screw extruder with twoshearing elements and one backflow element. Blending occurred at 190° C.with a rotor speed of 190 rpm. The melt was quenched upon exit from thedie in room temperature water, and was subsequently dried under vacuum.Film samples were compression molded at 190° C. for five minutes, andwere then quenched in room temperature water. The resulting film wasessentially amorphous.

Example 9 Extrusion-Cast Polylactic Acid Films

Polylactic acid (Natureworks 4041D) was exposed to blending in acounter-rotating twin screw extruder with two shearing elements and onebackflow element. Blending occurred at 190° C. with a rotor speed of 190rpm. The melt was quenched upon exit from the die in room temperaturewater, and was subsequently dried under vacuum. The processed pelletswere fed into a 1½ inch single screw extruder, were processed at 190°C., and were cast into film onto a chill roll kept at 36° C. that was 1cm away from the die exit. The resulting film was essentially amorphouswith zero orientation.

Example 10 Extrusion-Cast Films of Polylactic Acid with 3% by VolumeSubstituted Montmorillonite

Polylactic acid (Natureworks 4041D) was tumble mixed with 3% by volumeof substituted Montmorillonite (Nanocor I34). The dry blend was exposedto blending in a counter-rotating twin screw extruder with two shearingelements and one backflow element. Blending occurred at 190° C. with arotor speed of 190 rpm. The melt was quenched upon exit from the die inroom temperature water, and was subsequently dried under vacuum. Theprocessed pellets were fed into a 1½ inch single screw extruder, wereprocessed at 190° C., and were cast into film onto a chill roll kept at36° C. that was 1 cm away from the die exit. The resulting film wasessentially amorphous with zero orientation.

Film samples were cut into dumbbell shapes with width 46 mm, length 36mm, and thicknesses of approximately 0.5 mm. Cutting was performedaround a sample mold with a rotary blade at high speed to ensureintegrity of the sample edges. The machine used to follow the on-linetrue mechano-optical behavior is described in greater detail in Told, etal. (Toki, S., Valladares, D., Sen, T. Z., and Cakmak, M. “Real timebirefringence development of orientation in polymers during uniaxialstretching as detected by robust spectral birefringence technique”Annual Technical Conference—Society of Plastics Engineers (2001), 59th(Vol. 2), 1830-1834). The films was stretched to stretch ratios up to 5at various rates from 5%/min up to 3000%/min at 75, 85, and 95° C. forPLA and 105, 115, and 125° C. for PS.

FIG. 1 shows tire true stress—true strain behavior for compressionmolded PLA poly(lactic acid) stretched uniaxially to SR 5 at 75° C. and5%/min at clay loadings of 0, 1, and 3%. PLA with 0% clay loading doesnot reach the strain hardening point until near the very end of thedeformation, and reaches a true strain value of nearly 4. This is almostlike a taffy-pull behavior in which the bulk is deformed at such a lowrate that the chains orient and relax at about the same rate, leading toa high degree of thinning of the bulk at the center-point where themeasurements are made. However, with the addition of clay, thedeformation behavior changes significantly, leading to strain hardeningat much lower true strains, around 1.5 for PLA with 1% clay loading and1 for 3% clay loading. It is believed that the clay sheets act asinfinitely long, stiff chains, feeling the force more than thesurrounding bulk, and orienting in the direction of deformation muchmore easily. Due to their rigidity, the clay sheets cannot relax as thepolymer chains do, and they hold their orientation during deformation.Here, orientation is taking place in the absence of relaxation. Thepolymer chains entangle on and around the clay sheets will not beallowed to relax, as the clay cannot relax, and the polymer chains thathave a hydrogen bonding interaction with the clay sheets will alsoorient in the direction of deformation without relaxing. In the absenceof the relaxation mechanism, it is clear that strain hardening will takeplace at lower true strains. This preserves the width at the center ofthe sample, while creating more even deformation and properties in thematerial.

FIG. 2 shows the true stress—true strain behavior for PLA blended withclay loadings of 0, 1, 3, and 10% by volume, stretched at 50%/min at 75°C. to SR 5 and then held in constrained engineering strain for 30minutes at the stretching temperature. These results clearly illustratethe same general pattern of the results for the same films stretched at5%/min (with the exception of 10% clay loading, which was difficult tostretch due to its extreme brittleness). In this case, the film isstretched at such a rate that strain hardening clearly occurs at 0%loading; however, strain hardening again occurs at lower true strains asclay content is increased. Under these stretching conditions, strainhardening occurs at a true strain of approximately 2 for 0% loadedmaterial, and at 1.5 for 1% loading, 1 for 3% loading, and 0.5 for 10%loading. During constrained holding at 75° C. in length, true strainclearly relaxes in pure PLA film; however, the samples hold their finalshape in films with clay loading, and true strain does not relax.

FIG. 3 shows the plots of true mechanical behavior for PLA at clayloadings of 0, 1, 3, and 5% stretched at a rate of 500%/min and 75° C.to SR 5. Under these stretching conditions PLA loaded with nanoparticlesstrain hardens at a higher true strain that at lower rates, yet thetrend that strain hardening occurs at lower true strains with increasingclay content continues to hold, with PLA with 0% clay loading strainhardening at a true strain of approximately 1.75, 1% loaded material at1.65, 3% at 1.1, and 5% at 1. As for films stretched at 75° C. and50%/min, films with higher clay loadings stretched at 500%/min and 75°C. retain their final as-stretched shape to a higher degree with lessrelaxation of true strain.

The true mechanical behavior of PLA films with clay loadings of 0, 1,and 3%, stretched at 85° C. and 500%/min to SR 5 is shown in FIG. 4.Once again, strain hardening occurs at lower true strains with increasedclay loading. However, the effect of nanoparticles is significantlygreater at 85° C. than at 75° C. This is similar to what was seen in therate effect, with nanoparticle loading having a greater influence duringlower rate stretching.

FIG. 5 shows the true stress—true strain plots for PLA films loaded with0, 1, and 3%, nanoclay, stretched at 500%/min and 95° C. to SR 5. Underthese conditions, a taffy-pull behavior is strikingly present in theunfilled system. This stretching condition most clearly illustrates theimportance of nanoparticle loading in controlling strain hardening inpolymer systems, as strain hardening occurs at true strains less than orequal to 1.5 at a clay loading as small as 1% by volume. The ultimatetrue strains are 1.9 for 1% clay loading and 1.75 for 3%, while theunfilled system thins to a true strain approaching 4. There is a veryslight upturn from the rubbery plateau in the unfilled system at a truestrain of approximately 3; however, this flattens out.

FIG. 6 shows the true mechanical behavior of extrusion-cast films of PLAwith clay loadings of 0 and 3% stretched to SR 5 at 5%/min and 75° C.,then held in constrained length at 75° C. for one hour. Unlike thecompression-molded unfilled PLA film, the extrusion-cast unfilled PLAfilm does not show behavior reminiscent of a taffy-pull, and strainharden at a true strain of approximately 1.7. However, the filled systemalso shows improved strain hardening properties in this case, withstrain hardening occurring at 0.6 true strain. This is not only betterthan the unfilled system, but is 40% less than the correspondingcompression-molded system. The likely cause of this difference is thehigher orientation of the clay sheets in the as-cast extrudednanocomposite films in relation to the compression molded nanocompositefilms. During the casting process, the clay sheets orient in thedirection of the flow, and while the polymer chains relax to anunoriented state (the initial retardation of the film is zero), the claysheets act as though they have an infinite relaxation time due to theirhigh stiffness. When these films are subsequently stretched, thedeformation is taking place between the clay sheets, around whichpolymer chains are entangled and hydrogen bonded. This leads to higherlevels of stress, yet lower levels of birefringence, as some of thechains are locked in their place by the already oriented clay sheets.This behavior leads to the birefringence—true stress observations shownin FIG. 30. Unlike the results comparing compression-molded films, theStress Optical Rule behavior is different between the filled andunfilled systems, with the filled system possessing a lower StressOptical Constant. Also, the extrusion-cast nanocomposite shows earlierpositive deviation from linearity. This positive deviation into therelatively low second regime slope does indicate that strain-inducedcrystallization is hindered in the nanocomposite to some degree, though.This result gives further credence to the effect of “pre-oriented”nanoclay sheets, in that the crystallization is likely taking place inregions between the clay, sheets where the polymer chains may becomehighly oriented, as these are the areas that are likely sustaining thebulk of the deformation. The final level of birefringence is lower inthe case of the nanocomposite because the “pre-oriented” nanoclay sheetsblock significant deformation on a global level, whereas the unfilledsystem is free of these constraints.

During holding, true strain in the unfilled system rebounds drastically,while some true strain continues to develop in the nanocomposite. Thisis seen in the stress—optical data as a slight decrease in birefringenceduring holding for the unfilled system, and continued development ofbirefringence in the nanocomposite. While the “pre-oriented” clay sheetshindered chain motion during deformation, they also hinder chainrelaxation during holding. Holding at elevated temperature gives thechains time to crystallize from their limited orientation level, therebyresulting in continued development of birefringence and true strain.

In the birefringence—true stress behavior shown in FIG. 8, thenanocomposite film possesses a lower Stress Optical Constant than theunfilled PLA film. This is similar to what was seen in FIG. 30 for thesame films stretched at a lower rate; however, the difference betweenthe Stress Optical Constants is clearly greater during stretching at ahigher rate. Only the slightest positive deviation from Stress OpticalRule linearity occurs in the nanocomposite film, showing again thatstrain-induced crystallization is hindered by the presence of nanoclaysheets. Again, during holding, there is a significant decrease inbirefringence seen in the unfilled system, while the nanocompositeexperiences significant continued birefringence development during theholding stage, most likely due to crystallization.

As seen in the birefringence—true stress behavior of extrusion-caststretched at lower rates, FIG. 9 shows that there is a significantdifference between the initial Stress Optical Rule behavior of thefilled and unfilled systems, while continuing the trend that thisdifference becomes greater with increased stretching rate. While thereis a slight positive deviation from linearity in the unfilled system,the nanocomposite clearly shows negative deviation from Stress OpticalRule linearity, continuing the trend that nanoparticle loading resultsin hindered strain-induced crystallization and a higher tendency towardnon-Gaussian chain behavior. This higher tendency toward non-Gaussianchain behavior could either be due to the non-Gaussian segmentdistribution introduced by the nanoclay sheets acting as infinitelylong, infinitely stiff chains, or by the nanoclay sheets facilitatingchain orientation while acting as a steric or electrical polarizationbarrier to the formation of a crystalline phase.

The true mechanical data for extrusion-cast PLA and PLA nanocompositefilms stretched to SR 5 at 500% and at temperatures 85 and 95° C. isshown in FIG. 10. There is also a trend that as stretching temperatureincreases, the strain hardening point shifts to higher true strain. Thebirefringence—true stress data shows that as temperature increases,there is an increasing difference between the ultimate birefringencevalues of filled and unfilled systems. As in the case of films stretchedat lower rates, films stretched at higher temperatures produce moresimilarity between the initial linear portions of the curves for filledand unfilled systems. The trend also continues that the slope of thesecond regime is lower for the filled system than for the unfilledsystem, once again indicating that strain-induced crystallization issuppressed by the addition of nanoclay.

The true stress—true strain data for compression molded a-PS (atacticpolystyrene) films with clay loadings of 0, 1, and 3% by volume,stretched at 1000%/min and 115° C. are shown in FIG. 11. Although strainhardening is not as clearly seen in a-PS as it is in PLA, it is clearthat the addition of clay nanoparticles alone results in strainhardening. This very clearly illustrates that strain hardening isinduced by the addition of nanoparticles, as it does not even occur inthe unfilled system, while it clearly occurs in the films loaded withboth 1 and 3% by volume nanoclay. Due to the relatively large thickness(>1 mm) of these films and high rates of stretching, only a limitedamount of coherent birefringence data could be collected.

It is clear that the addition of nanoparticles improves the strainhardening of polymer films by lowering the true strain value at whichstrain hardening occurs. This will lead to films with higher integrityand more even properties. Not only can strain hardening in the rubberybe lowered significantly by the addition of nanoparticles, but thestrain hardening point can be easily controlled by slight changes in thecomposition of the blend, and great improvement can be achieved at theaddition of only 1% by volume. This not only allows for adaptation of anew material to an existing process, but that material will retain thelight weight of being a plastic, while gaining the end-use propertyenhancement that has been shown for nanocomposites in such processes asinjection molding.

Clearly, the process enhancement provided by nanoparticle loading is notlimited to slow crystallizing materials. The property enhancement seenin non-crystallizable a-PS shows that, in general, an essentiallyamorphous precursor system loaded with nanoparticles can have improvedstrain-hardening properties, when film is produced from this system andstretched in the rubbery state.

Extrusion-cast films perform better than compression-molded films due tothe “pre-orientation” of the nanoclay sheets induced by the process.This leads to earlier strain hardening, as the sheets will not changetheir orientation significantly during the early stages of deformation.The generalized mechanism of earlier strain hardening is presented inFIGS. 12A and 12B. The as-cast films is composed of essentiallyunoriented amorphous polymer chains and nanoclay sheets with varyingdegrees of intercalation and preferential orientation in the plane ofthe film (this preferential orientation in the plane of the film ishigher in extrusion-cast films). The clay sheets have favorableinteraction with the polymer chains, and form relatively strongtemporary network junctions with the chains while disturbing chain-chainentanglements. During deformation, the stiff clay sheets feel forcesmore than the flexible polymer chains, and quickly orient in thedirection of deformation. The high stiffness of the clay sheets leads tolonger relaxation times, and hinders the relaxation of oriented chains.This leads to earlier onset of strain hardening, as chains reach theirmaximal extension at lower true strains. This mechanism of increasingthe relaxation time is particularly effective at higher temperatures andlower rates, where chain relaxation typically has a greater effectduring deformation, as the clay sheets will have a much higherrelaxation time than the time of stretching regardless of rate, and areobviously not affected significantly by temperature in the range oftemperatures in which stretching is normally performed.

As can be appreciated, the process of the present invention can be useto produce a variety of polymeric objects and products, and so is notlimited to the examples of products which are provided. Clearly theproduction of films is benefited since they will result in uniformthicknesses and improved optical clarity. Such products might includephotographic film where the products achieve improved optical clarityand dimensional stability, including an ability of the film to uncurl.But, the process could also be used to make recording tape, or polymersfor photonic applications, such as optical retarders, as well as threedimensional objects, such as bottles or containers which would benefitfrom being grain hardened, dimensional stability, uniform thickness,optical clarity, and/or gas barrier characteristics.

The foregoing embodiments of the present invention have been presentedfor the purposes of illustration and description. These descriptions andembodiments are not intended to be exhaustive or to limit the inventionto the precise form disclosed, and obviously many modifications andvariations are possible in light of the above disclosure. Theembodiments were chosen and described in order to best explain theprinciple of the invention and its practical applications to therebyenable others skilled in the art to best utilize the invention in itsvarious embodiments and with various modifications as are suited to theparticular use contemplated. It is intended that the invention bedefined by the following claims.

1-10. (canceled)
 11. A process for controlling the strain hardeningproperties of a polymer comprising: blending a polymer and nanoparticlesto produce a polymeric composition; quenching the polymeric compositionat a temperature below the glass transition temperature of the polymerto yield an amorphous polymeric composition; forming a film from theamorphous polymeric composition; and subjecting the film to strainhardening in a rubbery state by stretching the film at a temperaturebetween the glass transition temperature and the cold crystallizationtemperature, wherein the nanoparticles are present in an effectiveamount of between 0.01% and 10% by volume based upon the volume ofpolymer used to form the polymeric composition in order to reduce thetrue strain at which the film formed from the polymeric compositionundergoes strain hardening, and wherein the steps of quenching thepolymeric composition and the step of forming the film areinterchangeable.
 12. The process of claim 11, wherein the polymer isselected from one or more homopolymers and copolymers of polyolefins,polyamides, polyimides, polyesters, aliphatic polymers, amorphouspolymers, crystallizing polymers, and blends, alloys and combinations oftwo or more thereof.
 13. The process of claim 11, wherein thenanoparticles are particles with at least one dimension in the nanoscaleselected from spheres, particles of irregular geometry, sheets, foils,fibers, wires, tubes or combinations of two or more thereof.
 14. Theprocess of claim 11, wherein the nanoparticles are selected from carbonnanoparticles, graphite nanoparticles, carbon nanotubes, graphitenanotubes, spherical nanoparticles, Buckyballs, glassy nanoparticles,silica-based nanoparticles, nanoclays, substituted Montmorillonite,metal oxide nanoparticles, metal sulfide nanoparticles, metal nitridenanoparticles, metal complex nanoparticles, metal nanoparticles,metallic alloy nanoparticles, metallic alloy nanowires, metallic alloynanospheres, metallic alloy nano-sized sheets, metallic alloy foils,colloidal nanoparticles, and mixtures of two or more thereof.
 15. Theprocess of claim 11, wherein the nanoparticles are substitutedMontmorillonite.
 16. The process of claim 11, wherein the nanoparticlesare present in an amount of between 0.1% and 10% by volume based uponthe volume of polymer used to form the polymeric composition.
 17. Theprocess of claim 11, wherein the nanoparticles are present in an amountof between 1% and 10% by volume based upon the volume of polymer used toform the polymeric composition.
 18. The process of claim 11, wherein thenanoparticles are present in an amount of less than 5% by volume basedupon the volume of polymer used to form the polymeric composition. 19.The process of claim 11, wherein the polymer composition is partially orcompletely molten when subjected to strain hardening.
 20. A strainhardened polymeric product produced from the polymeric composition ofclaim 11.